SC I ENCE ADVANCES | R E S EARCH ART I C L E

ELECTROCHEM ICAL ENERGY

1Department of Materials Science and Engineering, Stanford University, Stanford, CA94305, USA. 2Stanford Institute for Materials and Energy Sciences, SLAC National Ac-celerator Laboratory, 2575 Sand Hill Road, Menlo Park, CA 94025, USA.*Corresponding author. Email: yicui@stanford.edu

Wang et al., Sci. Adv. 2017;3 : e1701301 8 September 2017

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Ultrahigh–current density anodes with interconnectedLi metal reservoir through overlithiation of mesoporousAlF3 frameworkHansen Wang,1 Dingchang Lin,1 Yayuan Liu,1 Yuzhang Li,1 Yi Cui1,2*

Lithium (Li) metal is the ultimate solution for next-generation high–energy density batteries but is plagued from com-mercialization by infinite relative volume change, low Coulombic efficiency due to side reactions, and safety issuescaused by dendrite growth. These hazardous issues are further aggravated under high current densities needed bythe increasing demand for fast charging/discharging. We report a one-step fabricated Li/Al4Li9-LiF nanocomposite(LAFN) through an “overlithiation” process of a mesoporous AlF3 framework, which can simultaneously mitigatethe abovementioned problems. Reaction-produced Al4Li9-LiF nanoparticles serve as the ideal skeleton for Li metalinfusion, helping to achieve a near-zero volume change during stripping/plating and suppressed dendrite growth.As a result, the LAFN electrode is capable of working properly under an ultrahigh current density of 20 mA cm−2 insymmetric cells and manifests highly improved rate capability with increased Coulombic efficiency in full cells. Thesimple fabrication process and its remarkable electrochemical performances enable LAFN to be a promising anodecandidate for next-generation lithium metal batteries.

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INTRODUCTIONThe increasing demand from consumer electronics, electrical transpor-tation, and grid-scale storage calls for higher energy density and lower-cost batteries (1, 2). Despite the great commercial success of lithium-ion(Li-ion) batteries, the current battery chemistries based on graphiteanodes and transitionmetal oxide cathodes cannot meet the require-ments in the long run (3). Li metal, with the highest theoretical capacity(3860mA∙hour g−1) and the lowest potential (−3.040 V versus standardhydrogen electrode), is an attractive choice as the ultimate anode.More-over, it is an indispensable component in the next-generation Li-O2 andLi-S batteries (3–7). However, critical problems of Li metal anodes re-main unsolved. Specifically, a Limetal anode does not have a host con-fining Li inside, causing marked dimension change during Li stripping/plating (8, 9). The fragile solid electrolyte interphase (SEI) cannot with-stand the resulting mechanical deformation (10), such that it keepsbreaking and repairing throughout cycling, leading to thick SEI and rapidcapacity fading. Meanwhile, the cracks in SEI also locally concentrate theLi-ion flux, causing uncontrollable dendrite growth (11) and consequentsafety issues.Under increasedcurrentdensities for fast charging/discharging,accelerated volume fluctuation aggravates all the consequent hazardousissues, resulting in more severe battery failures.

Massive efforts have been made to solve the challenges of Li metalanodes, focusing onmodifications of electrolytes (12–17), SEIs (8, 18–24),and separators (25–27). Meanwhile, elaborative engineering of the anodestructure itself is also an important approach. Previous work includes useof Li powder instead of Li foil (10) and the fabrication of an exquisitecurrent collector with rigorously designed structures for Li metal deposi-tion (28–30). Although these different attempts effectively pushed forwardLi metal anode engineering technique and alleviated certain aspects of theproblems, few of them were able to simultaneously maintain a constantelectrode volume and eliminate dendrite growth. Recently, we proposed

the concept of stable Li metal host, realized by embedding Li metal intohollow carbon spheres (31), the space between layers of graphene oxides(9), polymer fiber matrices (32), and macroporous structures (33, 34).These modifications successfully reduced volume change, suppresseddendrite growth, and achieved better cycle performance. However, theseworks are limited to current densities below 5 mA cm−2. A higherworking current density would cause capacity fading and instability ofelectrodes. Nevertheless, fast charging/discharging (current densitygreater than 10 mA cm−2) is indispensable for advanced batteries. Thus,it remains a formidable challenge to design a Li metal anode capable ofoperating under these high current densities.

To realize a normal operation under ultrahigh current densities, webelieve that a Li metal anode would need to have the following prop-erties: (i) an ultrastable host accommodatingmetallic Li to eliminate thedamage on electrode and interfaces due to fast volume change duringhigh–current density operation (essentially the whole electrode needs tohave an almost zero volume fluctuation); (ii) a three-dimensional Limetal with a large electroactive surface area to effectively reduce thelocal current density; and (iii) an electrochemically inert,mechanicallystable and highly Li-ion conducting interphase to protect the structure.Here, we developed a novel Li metal anode, which simultaneously hasall the required properties. Using Al4Li9-LiF as a stable host, three-dimensional Li metal can be embedded into this nanocomposite (Li/Al4Li9-LiF) (LAFN) through a simple one-step “overlithiation” processof mesoporous AlF3.We demonstrated that LAFN is an ideal host withnear-zero volume change during stripping/plating. In addition, de-spite the space occupied by the skeleton, the metallic Li in the LAFNcomposite still offered a high specific capacity of ~1571 mA∙hour g−1

and a volumetric capacity of ~1447 mA∙hour cm−3. Meanwhile, LiFformed during overlithiation serves as a protectivematerial both electro-chemically and mechanically to improve the stability of the Li metalanode. Moreover, it also promotes the homogeneity of Li-ion diffusion,which can effectively suppress dendrite growth. Because of these supe-riorities, an LAFNanode is capable of achieving a long cycle life with lowoverpotential and high capacity and is able to work stably under an un-precedented high current density of 20 mA cm−2 for at least 100 cycles.

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RESULTSFabrication and morphology of LAFNSynthesis ofLAFNstarts fromLimetal foil andAlF3powder. Figure1 (D toF) shows the scanning electron microscopy (SEM) morphology ofthe pristine AlF3 powder. It is composed of secondary particles with adiameter of ~20 mm(Fig. 1D). Eachparticle consists of layers of columnarAlF3 arrays (Fig. 1, A, E, and F). These AlF3 arrays, with a column di-ameter of ~40 nm and spacing of ~20 nm, provide a mesoporousframework for Li reaction and infusion, making the AlF3 an ideal precur-sor for overlithiation. Hysteresis on the N2 adsorption-desorption iso-therm of AlF3 powder (fig. S1) further proves the mesoporous structure.Overlithiation of this mesoporous AlF3 framework follows reaction 1,where the alloy product is reasonably deduced through a Li-Al phase di-agram (35). Existence of the productswill be further proven through x-raydiffraction (XRD) characterization below.As shown in the equation, over-stoichiometric Li metal reacts with mesoporous AlF3 framework. In thisway, stoichiometricLimetal is able to reactwithAlF3, directlyparticipatinginto the framework construction, whereas excess Li metal fills into theinterspace of this framework. According to the reaction equation, massratio betweenAlF3 andLi in a stoichiometric reaction should be 0.4:0.17 g.Our overlithiation reaction used an overstoichiometric Li of 0.4:0.37 g.

Wang et al., Sci. Adv. 2017;3 : e1701301 8 September 2017

Theoretically, this ratio would give a final capacity of 1003 mA∙hour g−1

coming from the excess Li metal embedded inside.

4AlF3 þ ð21þ xÞLi→ Al4Li9 þ 12LiFþ xLiðexcessÞ ð1Þ

After puttingAlF3 andLi foil into a crucible, an elevated temperature of~350°C and continuous stirring were applied. After Li metal melted, itcould be separated into small droplets among the white powder of AlF3(fig. S2A). This phenomenon indicated the lithiophilic feature of themesoporous AlF3 framework; otherwise, lithium would become a largedroplet to reduce surface area. With continuous stirring, smaller andsmaller Li droplets were observed, whereas the color of AlF3 powdergradually turned from white to gray and then to black (fig. S2, B andC), indicating the formation of an Al4Li9-LiF skeleton (Fig. 1B). After aproper time of stirring, the remaining tiny lithium droplets would per-meate into the black powder of the Al4Li9-LiF skeleton, suggesting theinfusion of excess lithium into the framework (Fig. 1C). After the reac-tion, black powders of various sizes were obtained (fig. S2D). Later on,LAFNpowders were size-selected to remove very large particles (fig. S2E)and mechanically pressed into electrodes (fig. S2F). Electrode thicknesscould be controlled for various electrochemical testing. The pore spaces

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Fig. 1. Morphology of AlF3 powder and LAFN. (A) Sketch of AlF3 powder mesoporous framework. (B) Sketch of Al4Li9-LiF skeleton. (C) Sketch of LAFN. (D to F) SEMimages of AlF3 powder under different magnifications. Scale bars, 20 mm (D), 1 mm (E), and 200 nm (F). (G to I) SEM images of pristine LAFN surface, LAFN surface afterstripping 5 mA·hour cm−2 of Li metal, and LAFN surface after stripping and plating back 5 mA·hour cm−2 of Li metal. Scale bars, 2 mm. (J to L) Magnified SEM images of(G) to (I). Scale bars, 200 nm.

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between LAFN secondary particles were still maintained, and the electro-lyte was found to be very compatible with LAFN electrode (fig. S3). Asa result, Li ions can be transferred into the framework and used whereverneeded during cycling.

Figure 1 (G and J) shows the SEM images of LAFN powder underdifferent magnifications. After the overlithiation process, the columnararrays of pristine AlF3 were converted into particles about 50 to 100 nmin diameter because of the reaction with Li metal. Meanwhile, the po-rous structure was maintained after the reaction, with excess Li metalinfused into the host skeleton. Figure S4A shows a transmission electronmicroscopy image of a LAFN secondary particle. It can be observed thatthe secondary particle is composed of numerous small particles, inaccordance with the SEMmorphology. Energy-dispersive x-ray spec-troscopy (EDX) mapping was done for Al and F (fig. S4, B to D). Altends to concentrate in the center of the particle, whereas F spreadsout in the whole area, especially on the surface of the particle. Notethat the sporadic Al element observed in the outer part of the particlecan be attributed to the unreacted AlF3. Figure 1 (H and K) showsthe SEM images of LAFN morphology after stripping 5 mA∙hour cm−2

of metallic Li. The porous skeleton formed by Al4Li9-LiF nanopar-ticles could be obviously observed after Li stripping on the surface.After plating back 5 mA∙hour cm−2 of Li, most of the metallic Lideposited back into the porous structure, with Al4Li9-LiF nanopar-ticles faintly visible underneath (Fig. 1, I and L). This Al4Li9-LiF nano-particle framework with Li metal infusion structure can be furtherexhibited through sketch in Fig. 1C, in which the blue nanospheresrepresent the as-formed LiF nanoparticles, and the purple nano-spheres represent the Al4Li9 nanoparticles. These nanoparticles stackup to form a skeleton with interconnected pathways, filled with excessmetallic Li (transparent). From the reaction stoichiometry, we know thatthe ratio of Al to F should be 1:3, reasonably indicating a much largeramount of LiF particles than Al4Li9. In addition, this is also in accord-ance with the abovementioned EDX result, where Al4Li9 mostly concen-trates in the center of secondary particles and is surrounded by LiFnanoparticles. This structure will be further proven later.

Li dendrite suppression and near-zero volume changeMorphologies of Li foil and LAFN after different symmetric cell cyclesand current densities in carbonate-based electrolyte (see Materials andMethods) were observed. After one cycle of Li stripping and plating un-der a current density of 1 mA cm−2 and a capacity of 1 mA∙hour cm−2,the Li foil surface was covered with numerous dendrites, tangling witheach other (Fig. 2B). Meanwhile for the LAFN electrode, a smooth sur-face was maintained because of the confining effect and the homoge-neous growth among the framework (Fig. 2C). Low-magnificationSEM images can be seen in fig. S5 (A and B). After 10 cycles under acurrent density of 1 mA cm−2 and a capacity of 1 mA∙hour cm−2, thedendrite on the Li foil surface became much thinner, which could in-crease the possibility of piercing the separator, causing a soft short cir-cuit (fig. S6, A and C). However, for the LAFN electrode, no dendritegrowth was observed on the surface (fig. S6, B and D). Even when weincreased the applied current density to 20 mA cm−2, with capacityfixed at 1 mA∙hour cm−2, the LAFN electrode showed no Li metaldendrites (Fig. 2E) after one cycle of stripping and plating, whereasthe dendrites formed on Li foil surface curled up because of theincreased heterogeneity under higher current density (Fig. 2D). Low-magnification SEM images can be seen in fig. S5 (C andD). In addition,fig. S7 shows themorphology of LAFN electrodes after 50 and 100 cyclesunder a current density of 1mA cm−2 and a capacity of 1mA∙hour cm−2,

Wang et al., Sci. Adv. 2017;3 : e1701301 8 September 2017

and fig. S8 shows that of LAFN electrodes after 10 and 30 cycles under acurrent density of 1 mA cm−2 and a capacity of 3 mA∙hour cm−2. Bothfigures further prove the exciting capability of LAFN electrodes tosuppress dendrite growth even after long cycling. Furthermore, fig. S9shows the morphology of LAFN electrode (without any prestripping)after directly plating 1mA∙hour cm−2 and 5mA∙hour cm−2 of Limetalonto it. From the uniform surface of these electrodes, we can concludethat LAFN electrodes did not only offer a skeleton for Li metal but alsomodified interface chemistry for Li metal plating homogeneity.

Dendrite growth is successfully suppressedowing to three vital factors.The distinctive skeleton ofAl4Li9 surrounded by LiF nanoparticles playsa very important role. As shown in Fig. 2A, if we have a very conductiveskeleton, electrons can efficiently transfer to the anode surface, and thenthe Limetal will be plated back onto the surface of this electrode becauseof the high availability of both electrons and ions, leaving the skeletonreservoir empty. Dendrite growth will also be promoted because of theuneven surface and locally concentrated Li-ion flux. However, forLAFN, the insulating LiF nanoparticles (blue) dominate the skeletonand holdAl4Li9 (purple) inside (Fig. 2A), effectively restraining the elec-tron transfer to the surface. In this case, the framework is not very con-ductive. This promotes Li metal to be plated back inside the structureonto the underlying reserved Li, where electrons are available, graduallyfilling the whole skeleton. Then, dendrite growth can be suppressed to agreat extent because all stripped Li metals are used to fill the skeletonagain, and the nanocomposite structure can be effectively maintainedduring recurrent cycles. Besides, the effect of LiF itself suppressing den-drite growth has also been previously reported and proven (36–41).Li halide salts are capable of reducing the activation energy barrier forLi-ion diffusion at an electrode/electrolyte interface (14, 42). Thisincreased Li-ion diffusivity further promotes Li ion to transfer insidethe skeleton and plate onto existing Li metal instead of growing intodendrites on the surface. In addition, for Li foil, only surface Li hasaccess to electrolyte and Li cations, whereas LAFN reduces the effectivelocal current density and the Li cation flux around the nuclei due to theincreased surface area, thus increasing the deposition homogeneity.

Volume stability of LAFN was also characterized through thicknesstesting of electrodes after stripping and plating metallic Li. LAFN waspressed into the electrode, the cross-sectional image of which is shownin Fig. 2F. We made an ultrathick electrode of about 760 mm, aimingat increasing the volume change effect for better observation. Figure 2Gshows the cross-sectional image of an LAFN electrode after stripping37.5 mA∙hour cm−2 of Li metal inside (equivalent to ~181.93-mm-thickLi foil). The thickness shrinkage was only ~10 mm (~1.3% change).Essentially, the volume change was near zero, indicating that theelectrode is capable of maintaining the skeleton structure after Li metalstripping. Later, 37.5 mA∙hour cm−2 of Li metal was plated back, andFig. 2H shows that the electrode maintained a thickness of ~760 mmvery well. To guarantee the accuracy of this experiment, all threesampleswere cut from the same piece of LAFN electrode. The near-zerovolume change of LAFN is remarkable, for it not only indicates thatLAFN successfully provides a stable interconnected reservoir for Limetal for effective stripping and plating but also further confirms thecapability of Li metal plating back into the skeleton instead of grow-ing on the electrode surface, guaranteeing the structure reusabilityduring later cycling. Compared with the previous work of grapheneoxide host with ~20% volume change (9) and polymer fiber matrixwith 2.4% volume change (32), the near-zero volume change here isvery significant and will serve as a crucial factor for the cycling stabilityunder ultrahigh current densities.

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LiF-protected skeleton with Li-ion conductivityBasic electrochemical properties were further characterized for LAFN.Electrochemical impedance spectroscopy (EIS) was tested for as-assembled symmetric cells. Figure 3A shows the Nyquist plot forLAFN symmetric cells for the first 20 hours, and Fig. 3B shows thatfor Li foil. The high-frequency semicircles are good indicators of thecharge transfer resistance on the electrode-electrolyte interfaces. Be-cause of the unavoidable native oxide layer, Li foil showed an extremelyhigh interfacial resistance of ~1600 ohms, whereas for LAFN, a highlyreduced resistance of ~8 ohms could be observed. Impedance perform-ances were also tested after cycling of LAFN and Li foil symmetriccells (fig. S10). For Li foil, the charge transfer resistance rapidly de-creased to ~35 ohms mainly because of the broken native oxide layersas well as largely increased surface area caused by dendrite growth.Meanwhile, LAFN showed a rather stable and consistently lowresistance even after 100 cycles. This low interfacial resistance offers

Wang et al., Sci. Adv. 2017;3 : e1701301 8 September 2017

a further improved charge transfer capability compared to grapheneoxide (~30 ohms) (9). For Li foil, in contrast, continuous side reac-tions happen because of the exposure of Li to electrolyte, and the accu-mulated SEI layer with poor Li-ion conductivity leads to high resistance.These results guarantee that Li metal inside LAFN can be cycled withlower driving force for charge transferring. Furthermore, after strippingLi metal from one electrode to the other, charge transfer resistance ofLAFNsymmetric cells increased very slightly (fig. S11). Evenwhen 50%ofthe capacity (up to 30 mA∙hour cm−2) was stripped from one electrode,the charge transfer resistance only increased from ~8 to ~13 ohms. Thisfurther indicates the stable interface properties of LAFN during cyclingdue to its interconnected matrix.

Although an Al4Li9-LiF skeleton occupied a large part of the nano-composite, a high specific capacity could still be retained.After strippingLi metal inside LAFN electrode to 1 V versus Li+/Li under a very lowcurrent density of ~0.2 mA cm−2, the voltage profile (Fig. 3C) showed

Fig. 2. Dendrite-suppressed, volume-stable LAFN electrode. (A) Sketch of ideal and nonideal plating processes due to different conductivities of the skeleton. (B andC) SEM images of Li foil surface (B) and LAFN surface (C) after one symmetric cell cycle under 1 mA cm−2 and 1 mA·hour cm−2. Scale bars, 5 mm. (D and E) SEM images ofLi foil surface (D) and LAFN surface (E) after one symmetric cell cycle under 20 mA cm−2 and 1 mA·hour cm−2. Scale bars, 5 mm. (F to H) Cross-sectional SEM images of apristine LAFN electrode (F), an LAFN electrode after stripping 37.5 mA·hour cm−2 of Li (G), and an electrode after stripping and plating back 37.5 mA·hour cm−2 of Li (H). Scalebars, 500 mm.

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a total specific capacity of 1571 mA∙hour g−1. No second plateau canbe observed, indicating that only excess Li metal contributed to thiscapacity. This result is evenhigher than the theoretical capacity of excessLi metal (1003mA∙hour g−1), which can be attributed to the incompletereaction between AlF3 and Li. XRD was also applied to confirm this. InFig. 3D, the green curve of pristine LAFN confirms the formation ofAl4Li9-LiF framework and the existence of excess Li metal as well asAlF3 that have not completely reacted with Li metal (small peaks be-tween 50° and 60°). In addition, the blue curve shows the sample struc-ture after stripping Li metal to 1 V versus Li+/Li. The amorphous haloon the left side of these curves can be attributed to the Kapton tape forsealing. Compared to pristine LAFN, delithiated sample no longer had aLi peak, while still maintaining a high Al4Li9 peak. This feature can beobservedmore easily in amagnified figure (fig. S12), in accordance withthe delithiation voltage profile with no second plateau. X-ray photo-electron spectroscopy (XPS) was also applied for pristine LAFN and de-lithiated LAFN to further confirm its structure (fig. S13). The lack of Alpeaks on the surface of both pristine and fully stripped samples provesthat Al4Li9 alloy is in good protection of LiF nanoparticles. After de-lithiation, metal Li peak also disappeared, confirming the full strippingof excess Limetal. From these results, wemay reasonably deduce that allexcess Li was used during the delithiation process, and the Al4Li9framework was protected in good condition, possibly because of theprotection of electronically insulating LiF nanoparticles. The insulatingLiF confined Al4Li9 particles inside to maintain a stable skeleton whilestill offered an excellent interconnected pathway for Li metal and Liions. This feature further revealed the superiority of overlithiation aswell as the potential remarkable performance of LAFN battery systems.

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Electrochemical performance of LAFNHere, galvanostatic symmetric cell cycle capability was first com-pared between Li foil and LAFN. Figure S14 shows the voltage profileof Li foil and LAFN symmetric cells during the 1st, 10th, 50th, and100th cycles, with current density fixed at 1mA cm−2 and capacity fixedat 1 mA∙hour cm−2 in carbonate-based electrolytes. For the Li foil sym-metric cell, high overpotential could be seen at the beginning of chargingprocess during the first cycle, owing to the preobserved poor ionic con-ductivity of SEI on Li surface. High overpotential was also observed atthe beginning and end of discharging process, corresponding to theincreased difficulty of Li nucleation beneath thick SEI and the usageof previously unused fresh Li beneath the surface, respectively. This effectfaded with cycle number, because of the broken SEI layer on the surfaceof Li foil and the increased real surface area after dendrite growth.On thecontrary, LAFN showed flat, smooth voltage plateaus during chargingand discharging processes with a low overpotential below 40mV, as wellas better stability after 100 cycles. This superiority can be further provenin Fig. 4A, where LAFN symmetric cells offered a consistent, stable lowoverpotential about 35 mV during the first 100 cycles, much lower thanthat of Li foil counterparts, whose overpotential fluctuated between 60and 160 mV. Actually, long-term cycle capability was also tested, andthe result can be seen in fig. S15A. LAFNmaintained stable cycling after500 hours of testing, corresponding to 250 cycles, which could be seenfrom the enlarged figures from the 90th to 100th cycles (fig. S15B) and230th to 240th cycles (fig. S15C). Although the overpotential increasedfrom ~35 to ~45 mV, stable charging and discharging plateaus weremaintained. Meanwhile for Li foil, marked overpotential increase couldbe seen after 250 hours.

Fig. 3. Electrochemical properties of LAFN electrode. (A and B) Impedance performance of LAFN electrode (A) and Li foil electrode (B). (C) Voltage profile ofstripping Li from LAFN electrode to 1 V versus Li+/Li. (D) XRD characterization of LAFN electrode before and after stripping Li to 1 V versus Li+/Li. arb. unit, arbitrary unit.

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Later, cycling testswere carried out under increased current densitiesof 3 mA cm−2 (fig. S16), 5 mA cm−2 (fig. S17), 10 mA cm−2 (Fig. 4B),and 20 mA cm−2 (Fig. 4C), with capacity fixed at 1 mA∙hour cm−2.The overpotential of LAFN increased from 75mVunder 3mA cm−2,150 mV under 5 mA cm−2, to 240 mV under 10 mA cm−2, and thento 400 mV under 20 mA cm−2. Meanwhile, stable, smooth cyclingplateaus were maintained for the first 100 cycles. In contrast, theLi foil symmetric cell showed not only much higher overpotential butalso marked fluctuation during charging and discharging processes. Tobe more specific, cycling plateaus during different time periods undervarious current densities are directly plotted for a better comparison(fig. S18). It can be easily observed that the interconnected lithiumreservoir and faster kinetics of LAFN not only enabled it to work undermuchmitigated hysteresis, but also kept flat and smooth cycling plateauswithout fluctuations caused by local heterogeneity or soft short circuit.It should be noted that 20 mA cm−2 is an unprecedented high currentdensity that has rarely been used, if any, in the tests of Li metal anodes.For Li foil symmetric cells under 20mA cm−2, a high overpotential of~2 V was observed at the beginning of the first cycle (fig. S19), followedby a drop to a level even lower than that under 10mAcm−2, indicating asoft short circuit within the cells. In contrast, LAFN was capable ofworking stably under such a high current density for the first 100 cycles,with suppressed dendrite growth (Fig. 2E), verifying the potential ofmaintaining a stable interface and cycling performance under high cur-rent densities. In fig. S20, the voltage profile comparison is shown be-tween LAFN and Li foil symmetric cell cycling under metabolic current

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densities from0.5 to 10mAcm−2 andback to 0.5mAcm−2,with capacityfixed at 1 mA∙hour cm−2. Continuous flat charging and dischargingplateaus with obvious lower overpotential can be seen under all appliedcurrent densities for LAFN, manifesting its capability of working undermutative current densities for varied purposes. In addition, cyclingperformance under a current density of 1 mA cm−2, with areal capacitycorresponding to that of a commercial battery (3 mA∙hour cm−2), wasalso tested (fig. S21). Li foil symmetric cells showed a high overpotentialat the end of charging and discharging processes and stopped workingafter about 35 cycles. In comparison, LAFN symmetric cells worked nor-mally after over 400 hours (more than 70 cycles), with stable and lowvoltage plateaus in each cycle, manifesting potential of commercial levelcycle capability.

All these superiorities were further validated through full cell testing.Full cells using LAFNor Li foil pairedwith a LiCoO2 (LCO) cathodewerecompared through rate capability tests. Higher capacity was constantlyretained using LAFN as an anode, especially at higher rates. Here, weused an LCO cathode with an active material mass of ~3.6 mg. Usinga theoretical specific capacity of 140 mA∙hour g−1, we can calculatethat a rate of 1 C corresponds to a current density of 0.5 mA cm−2.The LCO/Li foil full cell only offered capacities of 125 mA∙hour g−1

under 0.2 C, 100 mA∙hour g−1 under 1 C, 73 mA∙hour g−1 under 4 C,and 5 mA∙hour g−1 under 10 C. In comparison, the LCO/LAFN fullcell gave capacities of 140mA∙hour g−1 under 0.2 C, 131mA∙hour g−1

under 1 C, 113 mA∙hour g−1 under 4 C, and ~80 mA∙hour g−1 under10 C (Fig. 5A). Here, all specific capacities were calculated on the

Fig. 4. Symmetric cell cycling performance of LAFN. (A to C) Symmetric cell cycling performance comparison between Li foil (red) and LAFN (blue) under 1 mA cm−2

(A), 10 mA cm−2 (B), and 20 mA cm−2 (C). Areal capacity was fixed at 1 mA·hour cm−2.

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basis of active LCOmass. LAFN can work properly under high rates,under which Li foil has already failed, further confirming the capabilityof LAFN working under high current densities. It should be noted thatthe capacity of the LCO/LAFN full cell also showedmuchbetter stabilitythan that of LCO/Li foil full cell. The capacity of the LCO/LAFN full cellunder a high rate of 10 C showed little fluctuation during the 10 cyclestested, with no trend of fading. Moreover, after resetting the rate to0.2 C, the LCO/LAFN full cell maintained the original capacity andwas very stable during the following 90 cycles, whereas the capacityof the LCO/Li foil full cell after resetting the rate back to 0.2 Cwas lowerthan that of the first cycle and kept decaying for the next 90 cycles.Cycling stability was also tested under 1 C with five cycles under0.2 C for electrode activation (fig. S22), and higher capacity with betterstability can still be observed for LAFN. These results in full cellsmay berelated to the reduced hysteresis using LAFN as an anode. Higheroverpotential could be observed for a Li foil anode at the beginningof each charging plateau (Fig. 5, B to D, dotted box, and fig. S23,A to C), owing to the sluggish kinetics at Li foil/electrolyte interface.This effect is also in accordance with symmetric cell cycling voltageprofiles. As we increased the rate, this hysteresis effect was also boosted(fig. S24, A to C), corresponding to the enhanced capacity gap betweenLCO/Li foil and LCO/LAFN full cells. The remarkable performance ofthe LAFN anode with the LCO cathode further verifies its potential asindustrialized Li metal anode material.

LAFN and Li foil were also paired with high–areal capacityLi4Ti5O12 (LTO; ~3.3 mA∙hour cm−2). LTO/LAFN cells still showeda highly increased rate capacity than LTO/Li foil cells, especially underhigher rates (fig. S25A). To be noted, here, a rate of 3 C correspondsto a current density as high as 10 mA cm−2. Thus, the capability ofLAFN working under high current densities is further verified. LAFNalso showed highly reduced hysteresis during full cell cycling, in accord-ance with LCO full cell results (fig. S25, B to E). Meanwhile, the mostimportant role of the high–areal capacity LTO cells was to serve as anindirect indicator ofCoulombic efficiency (CE) of the anodes. LTO itself

Wang et al., Sci. Adv. 2017;3 : e1701301 8 September 2017

does not provide Li during cycling, only serving as a Li reservoir. More-over, theCEof LTO is near 100%except at initial cycles. As a result, onlyanodes serve as Li sources, and the loss of Li can be mostly attributed tothe low CE of the anodes. We paired the high–areal capacity LTOelectrode with 50-mm-thick Li metal foil (~10 mA∙hour cm−2) andLAFNelectrodewith a capacity of ~16mA∙hour cm−2 (thinnest electrodecan be fabricated using LAFN powder) and tested their cycle life under0.5 C. The point where specific capacity begins to fade markedly in-dicates the cycle number used to consume excess Li metal in LAFN orLi foil. As shown in fig. S26, capacity of LTO/Li foil cell started to fadeat the 28th cycle, illustrating a CE of ~92.7%. In contrast, the capacityof LTO/LAFN cell remained almost constant even after 150 cycles,giving a CE of at least ~97.4%, showing further improvements evencompared to previously reported polymer fiber matrix (32).

DISCUSSIONLAFN as a Li metal anode was synthesized through the overlithiationprocess of a mesoporous AlF3 framework. LAFN had a porous skeletonwith an interconnected Li metal reservoir. This designing not onlyoffered a stable host for lithium metal during stripping and platingbut also took advantage of the reaction-produced insulating LiF forthe electrochemical protection of the Al4Li9 nanoparticles as well asLi cation diffusion promoter for increased Limetal plating homogeneity.These superiorities endowed LAFN with suppressed dendrite growthand near-zero volume change. As a result, LAFN showed much stablecharging and discharging plateaus in symmetric cells, much highercapacity capability in full cells using LCO/LTO cathodes, and increasedCE, compared to Li foil. The LAFN anode worked properly under anunprecedented high current density of 20mA cm−2, under which Li foilshorted immediately. Considering the safe and single-step fabricationprocedure, we believe that LAFNmay inspire designs of next-generationLi metal anode that can accommodate fast charging/dischargingdemands in the near future.

Fig. 5. Electrochemical performance of LCO/LAFN cells. (A) Rate capability of LCO/LAFN cells and LCO/Li foil cells at various rates from 0.2 to 10 C. (B to D) Voltageprofile comparison between LCO/LAFN cells and LCO/Li foil cells at rates of 1 C (B), 2 C (C), and 4 C (D). The dotted boxes are enlarged in fig. S23 (A to C).

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MATERIALS AND METHODSLAFN fabricationAlF3 powder (Aldrich) was transferred into an argon-filled glove boxwith sub–parts per million O2 level for later use. AlF3 powder (0.40 g)was firstly put into a stainless steel crucible on a hot plate heated up to~350°C.Afterward, 0.37 g of Li foil (99.9%;AlfaAesar)wasweighted, cutinto small pieces, and put into the crucible. Excess lithium foil ratio couldbe adjustedbutwithin 0.40 g to guarantee auniformpowder-like product.After the lithium foil melted, the AlF3 powder and the liquid Li metalwere vigorously stirred, and lithium would become increasingly smallerdroplets. After the white AlF3 powder turned to black and lithium smalldroplets started to wet the powder, heating was stopped and stirring wascontinued.After cooling down, LAFNpowderwas obtained. Tomake anLAFN electrode, the as-obtained LAFN powder was mechanicallypressed, using a pellet die (Aldrich) with ~10 metric tons of pressure.Electrode thickness could be controlled through manipulation of thepowder weight. All processes were done in a glove box.

CharacterizationsSEM images were taken with a FEI XL30 Sirion scanning electron mi-croscope. For the SEM characterizations on batteries after cycling, theelectrodes were first extracted from the coin cells in the glove box andfollowed by gentle rinse in 1,3-dioxolane (DOL) to remove Li saltresidue. XRD patterns were recorded on a PANalytical X’Pert instru-ment. To protect the highly reactive Li-related compound from the air,the XRD samples were loaded on a glass slide and covered with Kaptontape in the glove box before the XRD measurements. XPS analysis wasobtained on an SSI SProbe XPS spectrometer with an Al (Ka) source.

Electrochemical measurements on symmetric and full cellsTo study the Li stripping/plating processes within a symmetric cell con-figuration, 2032-type coin cells (MTI) were assembled here. The anodesused here were either the abovementioned LAFN electrodes or the Lifoils. To prepare LCO cathodes, lithium cobalt oxide power (MTI) wasmixedwith carbon black (TIMCAL) and polyvinylidene fluoride (MTI)with a mass ratio of 8:1:1, with N-methyl-2-pyrrolidone as the solvent.The mass loading for LCO cathodes was ~4.5 mg cm−2. To preparehigh–areal capacity LTO electrodes, LTO (MTI), polyvinylidene fluo-ride (MTI), and carbon black (TIMCAL) were mixed in a ratio of8:1:1, with N-methyl-2-pyrrolidone as the solvent. The areal massloading of LTO electrodes was ~24 mg cm−2. All electrodes used hadan area of ~1 cm2. The LAFN electrodes used for symmetric cell testingand LCO full cell testing were ~400 mm thick, offering an areal capacity~60 mA∙hour cm−2. The electrolyte used were 1 M lithium hexa-fluorophosphate (LiPF6) in 1:1 ethylene carbonate/diethyl carbonate(BASF Selectilyte LP40), with 10% fluoroethylene carbonate (NovolyteTechnologies Inc.) and 1% vinylene carbonate (Novolyte TechnologiesInc.) as additives. Celgard 2325 (25-mmpolypropylene/polyethylene/polypropylene) was used as the separator. The control bare Li cellswere assembled using freshly scraped Li foil. Galvanostatic cyclingwas conducted either on an Arbin 96-channel battery tester or on aLAND 8-channel battery tester. The EIS measurements were carriedout on a Biologic VMP3 system.

SUPPLEMENTARY MATERIALSSupplementary material for this article is available at http://advances.sciencemag.org/cgi/content/full/3/9/e1701301/DC1fig. S1. N2 adsorption-desorption isotherm of AlF3 powder.fig. S2. Fabrication of LAFN.

Wang et al., Sci. Adv. 2017;3 : e1701301 8 September 2017

fig. S3. Digital photo of 20 ml of ethylene carbonate/diethyl carbonate electrolyte droppedonto Li foil and LAFN electrode.fig. S4. Transmission electron microscopy characterization of pristine LAFN.fig. S5. Morphology of LAFN after symmetric cell cycling.fig. S6. Morphology of LAFN after symmetric cell cycling.fig. S7. Morphology of LAFN after symmetric cell cycling.fig. S8. Morphology of LAFN after symmetric cell cycling.fig. S9. Morphology of LAFN after directly plating Li metal onto its surface.fig. S10. EIS performance of Li foil and LAFN symmetric cells after different cycles.fig. S11. EIS performance of LAFN symmetric cells after stripping different percentages ofcapacity from one side to the other.fig. S12. Enlarged XRD profile of the dotted box in Fig. 3D.fig. S13. XPS of LAFN.fig. S14. Voltage profile of Li foil and LAFN symmetric cell during the 1st, 10th, 50th and 100thcycles, with a current density of 1 mA cm−2.fig. S15. Long-term cycling stability of LAFN electrode.fig. S16. Symmetric cell cycling performance of LAFN under a current density of 3 mA cm−2.fig. S17. Symmetric cell cycling performance of LAFN under a current density of 5 mA cm−2.fig. S18. Voltage plateau comparison.fig. S19. Symmetric cell cycling performance of LAFN under a current density of 20 mA cm−2.fig. S20. Symmetric cell cycling performance of LAFN under metabolic current densities.fig. S21. Symmetric cell cycling performance of LAFN under an areal capacity of 3 mA·hour cm−2.fig. S22. Electrochemical performance of LCO/LAFN cells.fig. S23. Enlarged voltage profile of LCO full cell batteries.fig. S24. Electrochemical performance of LCO/LAFN cells.fig. S25. Electrochemical performance of LTO/LAFN cells.fig. S26. Electrochemical performance of LTO/LAFN cells for CE testing.

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AcknowledgmentsFunding: Y.C. acknowledges the support from the Assistant Secretary for Energy Efficiencyand Renewable Energy, Office of Vehicle Technologies of the U.S. Department ofEnergy under the Battery Materials Research Program and Battery 500 Consortium. Authorcontributions: H.W. and Y.C. conceived the concept and designed the experiments.D.L. helped set experiment procedures. H.W. performed the materials synthesis. H.W., D.L.,Y. Liu, and Y. Li conducted the characterizations. H.W. performed the electrochemicalmeasurement. H.W. and D.L. fabricated the LCO and LTO full cell batteries. H.W. andY.C. co-wrote the paper. All authors discussed the results and commented on themanuscript. Competing interests: The authors declare that they have no competinginterests. Data and materials availability: All data needed to evaluate the conclusions inthe paper are present in the paper and/or the Supplementary Materials. Additional datarelated to this paper may be requested from the authors.

Submitted 23 April 2017Accepted 7 August 2017Published 8 September 201710.1126/sciadv.1701301

Citation: H. Wang, D. Lin, Y. Liu, Y. Li, Y. Cui, Ultrahigh–current density anodes withinterconnected Li metal reservoir through overlithiation of mesoporous AlF3 framework. Sci.Adv. 3, e1701301 (2017).

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framework3of mesoporous AlFcurrent density anodes with interconnected Li metal reservoir through overlithiation−Ultrahigh

Hansen Wang, Dingchang Lin, Yayuan Liu, Yuzhang Li and Yi Cui

DOI: 10.1126/sciadv.1701301 (9), e1701301.3Sci Adv 

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